Permanent magnet alloys for gap magnets

ABSTRACT

Provided are Ce/Co/Cu permanent magnet alloys containing certain refractory metals, such as Ta and/or Hf, and optionally Fe which represent economically more favorable alternative to Sm-based magnets with respect to both material and processing costs and which retain and/or improve magnetic characteristics useful for GAP MAGNET applications.

RELATED APPLICATION

This application claims benefit and priority of provisional applicationSer. No. 62/708,546 filed Dec. 12, 2017, the entire disclosure anddrawings of which are incorporated herein by reference.

CONTRACTUAL ORIGIN OF THE INVENTION

This invention was made with government support under Grant No.DE-AC02-07CH11358 awarded by the Department of Energy. The Governmenthas certain rights in the invention.

FIELD OF THE INVENTION

The present invention relates to permanent magnet alloys and topermanent magnets made from the alloys.

BACKGROUND OF THE INVENTION

There now exists a “GAP” (an empty niche of magnetic energy products)between present-day low-flux (ferrites, Alnico) and high-flux(Nd₂Fe₁₄B-type and SmCo₅-type) magnets.

For example, there is interest in so-called “GAP MAGNETS” that reachenergy products of about 10-20 MGOe and that would be both advantageouswith respect to the cost of constituents and their fabrication intopermanent magnets as well as reliable in performing successfully withinthe “GAP” which exists between present-day low-flux (ferrites, alnico)and high-flux (Nd₂Fe₁₄B- and SmCo₅-type) magnets.

Despite previous extensive explorations, the intrinsic properties of theCeCo₅-type systems have not been fully or systematically established,and the metallurgy related to the magnetic pinning/coercivity mechanismis not fully understood. Although anisotropy characterization is bestobtained from single crystals, single crystal growth reports in Cu or Fesubstituted CeCo₅ systems are scarce and limited to several Bridgemantype attempts [see G. Chin, et al., “Directional solidification ofCo—Cu—R permanent-magnet alloys,” IEEE Transactions on Magnetics 8,29-35 (1972) and see E. A. Nesbitt, et al. “Intrinsic magneticproperties and mechanism of magnetization of Co—Fe—Cu—R permanentmagnets,” in AIP Conference Proceedings (AIP, 1973)] in the vicinity ofthe composition about CeCo_(3.5)Fe₀₅Cu [see H. Okamoto, “Ce—Co PhaseDiagram, ASM Alloy Phase Diagrams Database, P. Villars, editor-in-chief;H. Okamoto and K. Cenzual, section editors,” 1990].

SUMMARY OF THE INVENTION

The present invention provides permanent magnet alloys comprising Ce,Co, and Cu also containing one or more of certain refractory metals(e.g. one or more of Ta, Hf, Zr, Nb, Mo, and W) and optionally Fe whichrepresent economically more favorable alternative to Sm-based magnetswith respect to both material and processing costs and can retain and/orprovide improved magnetic characteristics. The magnet alloys are usefulfor making so-called “GAP MAGNETS”.

The present invention envisions substitution of more than 50% ofcritical rare-earth elements, i.e., Sm, Dy, Nd etc, in the commercialhigh-flux permanent RCo₅ magnets by cheaper, more abundant andnon-critical Ce. This will significantly reduce costs of material,whereas the performance of such magnet must clearly surpass the levelsof known commercial non-rare earth grades, and/or may reach thecharacteristics of the best rare-earth containing representatives.

The present invention envisions a “GAP MAGNET” that utilizes widelyavailable and inexpensive Ce as more affordable alternative to criticalrare-earths, making the magnet significantly cheaper and less supplydependent, and yet successfully performing in the range of 10-20 MGOe;i.e., within the “GAP” (an empty niche of energy products) which existsbetween present-day low-flux (ferrites, alnico) and high-flux(Nd₂Fe₁₄B-type and SmCo₅-type) magnets.

In addition, the present invention envisions further reductions inmaterial costs by a partial substitution of Co by Fe. The substitutionFe was observed to strongly improve both the Curie temperature andmagnetization of the magnet at room temperature. This substitution willlower the heat treatment requirements, which consequently willfacilitate simplification of manufacturing process, i.e., decreasing thenumber of processing stages and reducing energy consumption.

Still further, the present invention envisions heat treating the abovepermanent magnet alloys in a manner to develop a bi-modal microstructurehaving refractory metal-containing laminar precipitates accompanied withnearby Cu-enriched and Co-depleted regions in a microstructure matrix,which can be primarily a single crystal grain and wherein formation ofthe bimodal microstructure produces microstructural changes thatcontribute to dramatic improvement of magnetic properties after heattreatment.

These and other advantages of the present invention will become morereadily apparent from the following detailed description taken with thefollowing drawings.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows PXRD pattern and Rietveld refinement results for EXAMPLE#1; i.e., Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) (crushed and powderedsingle crystals). The observed profile is indicated by circles and thecalculated profile by solid line. Bragg peak positions are indicated byvertical tics, and the difference is shown at the bottom. The insert ofFIG. 1 shows the temperature-time profile of the flux growth.

FIGS. 2a-2c show SEM analysis or as-grown single crystals of the EXAMPLE#1; i.e., Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy, which arebackscattered electron images of the [001] surface with differentmagnifications, FIG. 2d shows the target areas for elemental X-rayanalysis; i.e., target squares 1-3.

FIG. 3a shows magnetic hysteresis loop for as-grown single crystal ofthe EXAMPLE #1; i.e., Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy. FIG. 3bshows the demagnetization 4πM and B curves. FIG. 3c shows an estimationof the energy product (BH)_(max).

FIG. 4 shows a PXRD pattern and Rietveld refinement results for heattreated “COMPOSITE CRYSTAL” of EXAMPLE #1; i.e.,Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy. The observed profile isindicated by circles and the calculated profile by solid line. Braggpeak positions are indicated by vertical ties, and the difference isshown at the bottom. The insert of FIG. 4 shows the temperature-timeprofile of the heat treatment.

FIG. 5a-5d show SEM back scattered electron images of the heat treated“COMPOSITE CRYSTAL” of EXAMPLE #1; i.e.,Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy. FIGS. 5a-5d are backscatteredelectron images of the [001] surface with different magnificationsshowing bimodal laminar microstructure.

FIG. 6a shows a back scattered electron image of the [001] surface ofthe heat treated “COMPOSITE CRYSTAL” of EXAMPLE #1; i.e.,Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy, and the target areas for SEMelemental X-ray analysis; namely, the target squares 1-3 (matrix) andtargets 4 and 5 (laminas). Scale bar is 25 μm. A target area 6(general/total image) investigated the general matrix. FIG. 6b shows SEMX-ray results for the target areas 1, 2, 3, 4, 5, and 6.

FIG. 7a shows high resolution TEM images of as-grown crystals of EXAMPLE#1, i.e., Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy. FIG. 7b shows theTEM results for heat-treated crystals of EXAMPLE #1, i.e.,Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy.

FIG. 8a shows magnetic hysteresis loop of the heat treated “COMPOSITECRYSTAL” of EXAMPLE #1; i.e., Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy.FIG. 8b shows the demagnetization 4πM and B curves. FIG. 8c shows anestimation of the energy product (BH)_(max).

FIG. 9a shows the magnetic properties of “COMPOSITE CRYSTAL” incomparison to un-isotropic ingot of the EXAMPLE #1; i.e.Ce_(15.1)Ta_(0.6)Co_(67.8)Cu_(16.1) alloy wherein magnetic hysteresisloops of the “COMPOSITE CRYSTAL” (gray circles) of the “COMPOSITECRYSTAL” and of non-isotropic ingot (black circles) are shown. FIG. 9bshows the magnetic TGA measurement for the determination of the Curietemperature.

FIG. 10 shows the magnetic hysteresis loop of the bulk, un-alignedsamples containing Fe substituted for Co and of the Fe-free alloy ofEXAMPLE #1 showing improvement of both coercivities and magnetizationsin the Fe-substituted sample, i.e., H_(c) from about 6.3 to about 7.6kOe, B_(r) from about 2.1 to about 4.7 kG and M_(s) from about 3.2 to5.2 kG.

FIG. 11a shows the demagnetization 4πM and B curves of EXAMPLE #3 i.e.,the Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4) alloy, in comparison toEXAMPLE #1. FIG. 11b shows an estimation of the of the energy product(BH)_(max.) of EXAMPLE #3 in comparison to that of Example #1.

FIG. 12a shows a back scattered electron image of the [001] surface ofas-grown crystal of EXAMPLE #3; i.e.,Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4) alloy. Scale bar is 10 μm.FIG. 12b shows a back scattered electron image of the [001] surface of“COMPOSITE CRYSTAL” of heat-treated EXAMPLE #3; i.e.,Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4) alloy.

FIG. 13a-13c are SEM backscattered electron images of samples wherein(FIG. 13a )-sample III—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9); (FIG. 13b)-sample IV—Ce_(16.3)Ta_(0.3)Co_(61.7)Cu_(21.7); and (FIG. 13c )-sampleV—Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4), wherein the upper panelsare before and the lower panels are after heat treatment. All imageswere taken at a magnification 5000× and 15 kV.

FIG. 14a is a HAADF STEM image of as-grown sampleIII—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9) showing the overallmicrostructure. FIG. 14b is an enlarged HAADF image shows a darkcontrast line. FIG. 14c is a diffraction pattern taken from the regionshown m FIG. 14b including the matrix and the dark line. FIG. 14d is ahigh resolution STEM image taken from left boxed area in FIG. 14b under[1-10] zone axis. The inset at bottom right is an enlarged atomic imagewith atomic model of hexagonal 1:5 Ce/Co/Cu phase. The bright dots anddark dots in the images correspond to atomic columns of Ce and (Co, Cu)elements, respectively. FIG. 14e is an enlarged image of right boxedarea in 14b and dark line in single crystalline phase is shown clearly.FIG. 14f shows EDS elemental mapping of FIG. 14e indicating Coenrichment in the line, the small Co and Cu elemental maps-insets arepresented for contrasting observation of Cu depletion in the same line.

FIG. 15a is a HAADF image of the heat treated sample III showing theoverall microstructure. FIG. 15b is an enlarged image of the circledarea of FIG. 15a . FIG. 15c is an EDS elemental mapping corresponding toFIG. 15b wherein the bright regions are Ta-rich and considered as a Taprecipitate. FIG. 15d is an enlarged image of the boxed area of FIG. 15b. FIG. 15e is a diffraction pattern taken from FIG. 15d including thematrix and the Ta precipitate. FIG. 15f is a high resolution STEM imagetaken from the matrix in FIG. 15d under [1-0] zone axis. Scale bar is 2nm.

FIG. 16a is a high resolution HAADF image of the interface (dashedlines) between the matrix and the Ta precipitate taken from boxed areain FIG. 15d . FIGS. 16b-16f show corresponding EDS elemental mappingresults; FIG. 16b —Ta,Co,Ce, and Cu, FIG. 16c —Ta, FIG. 16d —Co, FIG.16e —Ce, FIG. 16f —Cu. The white dashed lines indicate the same positionin each image.

FIG. 17 shows Curie temperatures for the as-grown samples I, II, III,IV, and V inferred from the peaks in derivative of magnetization withrespect to temperature; i.e. dM/dT obtained for each crystal sample I-V.Magnetization data were obtained under magnetic field of 0.01 T.

FIG. 18 shows representative M(H) isotherms for the as-grown crystalsample V. In the inset, spontaneous magneetizations for each temperatureinferred from the extrapolation of the linear regions of M(H) back toH=0. Star shows extrapolated T_(c) value following the Bloch law:M(T)=M(0)(1−(T/T_(c))^(3/2)).

FIG. 19 shows a comparison of temperature dependent magnetocrystallineanisotropy energy density of sampleIII—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9) and sampleV—Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4).

FIG. 20 shows room temperature second quadrant magnetic hysteresis loopsfor the as-grown crystals sample III—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9)and sample V—Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4) where 4πM isindicated as solid line and B as a dashed line in the left panel.Estimation of the energy products (BH)_(max.) is shown in the rightpanel.

FIG. 21 shows magnetic hysteresis loops of the heat-treated crystalsamples I—Ce_(15.1)Ta_(1.0)Co_(74.4)Cu_(9.5), sampleII—Ce_(16.3)Ta_(0.6)Co_(68.9)Cu_(14.2), sampleIII—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9), sampleIV—Ce_(16.3)Ta_(0.3)Co_(61.7)Cu_(21.7), and sampleV—Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4) at 300 K. In FIG. 21, theupper-left inset shows dependence of the spontaneous magnetization M_(s)vs. Cu concentration for the as-grown and heat treated crystals. Thelower-right inset shows dependence of the coercivity H_(c) vs. Cuconcentration for the as-grown and heat-treated crystals.

FIG. 22 shows room temperature second quadrant magnetic hysteresis loopsfor the heat-treated crystal sampleIII—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9) and sampleV—Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4) where 4πM is indicated assolid line and B as a dashed line in the left panel. Estimation of theenergy products (BH)max. is shown in the right panel.

FIG. 23a shows a back scattered electron image of the polished surfaceof the heat treated polycrystalline sample of EXAMPLE #11; i.e.,Ce_(15.0)Hf_(0.7)Co_(61.8)Fe_(12.1)Cu_(10.4) alloy, and the target areasfor SEM elemental X-ray analysis; namely, the target squares 0 1-0 4(matrix) and targets 0 5 1-0 5 3 (laminar precipitates). Scale bar is 50μm. FIG. 23b shows SEM/EDS X-ray results for the target areas 1, 2, and0 b 1. FIG. 23c contains Table 8 which shows the elemental atomicfractions for target areas of matrix (0 1-0 4) and Hf-rich precipitates(0 5 1-0 5 3).

FIG. 24 shows the demagnetization 4πM and B curves of EXAMPLE #11 i.e.,the Ce_(15.0)Hf_(0.7)Co_(61.7)Fe_(12.1)Cu_(10.4) alloy in the left paneland shows an estimation of the of the energy product (BH)_(max) of theEXAMPLE #11 in the right panel.

DETAILED DESCRIPTION OF THE INVENTION

The present invention provides Ce/Co/Cu permanent magnet alloys whichcontain one or more of certain refractory metals that can include atleast one of Ta, Hf, Zr, Nb, Mo, and W, and optionally Fe and which canbe heat treated to promote a bimodal microstructure having refractorymetal-containing laminar precipitates accompanied with Cu-enriched andCo-depleted regions in the single crystal matrix microstructure. Suchpermanent magnet alloys comprise controlled amounts of Ce, Co, Cu andrefractory metal, and optionally Fe. An illustrative embodiment of theinvention involves a permanent magnet alloy that comprises, in atomic %,about 13 to about 17% Ce, about 38 to about 70% Co, about 10% to about30 atomic % Cu, optionally about 10% to about 20% Fe, wherein the alloyincludes at least one refractory metal in an individual amount orcollective amount (if more than one refractory metal is present) greaterthan 0, such as at least about 0.1 up to about 3 atomic % of the alloycomposition. Embodiments of the invention envision including two or moreof the above refractory metals, such as for example both Ta and Hf, inthe alloy composition to tailor microstructure and/or magneticproperties of the resulting magnet to particular applications.

A further illustrative embodiment of the invention involves a permanentmagnet alloy comprising, in atomic %, about 14 to about 15.5% Ce, about57.5 to about 62.0% Co, about 10% to about 16.5 atomic % Cu, about 10%to about 12.5% Fe, wherein the alloy includes at least one refractorymetal in an individual amount or collective amount (if more than onerefractory metal is present) of about 0.5 to about 1 atomic %. Forexample, when both Ta and Hf are present in the alloy, the collectiveamount thereof is about 0.5 to about 1 atomic %.

The magnet alloy is subjected to a series of particular solution heattreatments typically at solution temperatures 1000-1100° C. and agingprocedures to develop a bimodal laminar microstructure. For purposes ofillustration and not limitation, a series of heat treatments involveheating at solution temperatures of 1000-1100° C. (e.g. for 1-1.5 days)followed by aging at 400-450° C. (e.g. for 0.5-1 days). The heat treatedalloy modified pursuant to the invention can deliver magneticcharacteristics acceptable for the so-called “GAP MAGNET”; namely,T_(c)>300° C., H_(c)=0.5-1.0 T, R_(f)=4-8 kG and (B)_(max.)=7-15 MGOe.

The present invention will be described below with respect to thefollowing Examples that are offered for purposes of illustration and notlimitation with respect to the scope of the invention.

EXAMPLES

This EXAMPLE #1 illustrates an initial experiment which resulted inrealization of a Ce-substituted, Ta-doped RCo₅-type magnet pursuant toan illustrative embodiment of the invention which had a particular alloycomposition represented by Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1).

Well-formed plate-like crystals were obtained during self-flux singlecrystal growth from an initial loading compositionCe₁₈Co_(62.32)Cu_(19.68) in a Ta container at 1050-1070° C. The excessof flux was decanted by centrifuging at nearly reaction temperature. Theexact temperature profile of this crystal growth is presented in FIG. 1(see insert). Rietveld fitting of powder X-ray pattern taken fromcrushed single crystals, FIG. 1, shows that all Bragg reflections arewell indexed within CaCu₅-type structure (P6/mmm. a=4.943(1), c=4.032(1)Å), providing strong argument for single-phase nature of the as-growncrystals. Additionally, SEM backscattered electron images of thesecrystals, FIGS. 2a-2c , revealed uniformity of their polished [001]surface (even at magnifications close to ×30000) suggesting theirsingle-phase nature. The elemental SEM X-ray analysis, FIG. 2d and Table1 below, revealed that crystals represent, about 15-16 atomic % Cu-dopedCeCo₅ phase with minor but critical Ta content of about 0.5-0.6 atomic%.

TABLE 1 Atomic Fractions Area Co Cu Ce Ta 1 68.12 15.70 15.59 0.59 267.80 16.11 15.53 0.56 3 67.64 15.95 15.86 0.55

The minor presence of Ta, Table 1, is explained by slight dissolution ofinner wall surfaces of Ta reaction container and diffusion of Ta atomsinto reaction liquid during the long term dwelling process at themaximum temperature of 1200° C. in the high-temperature furnace for 9-10hours as well as at ramping down to 1050° C. for 75 hours. Since no Taprecipitation or/and segregation was revealed during both SEM and XRDanalyses, it was concluded that Ta was being either incorporated intothe crystals structure interstices or uniformly distributed in the formof nano-sized precipitates, the detection of which is beyond theresolution of both instrumental methods applied.

Interestingly, these single crystal samples showed magnetic hysteresiswhen measured along the easy axis of magnetization [001] with H_(c)=1.6kOe and B_(r)=4.2 kG, M_(s)=about 4.2 kG and (BH)_(max)=about 3.7 MGOe,FIGS. 3a-3c . This by itself is remarkable considering common beliefthat appearance of the coercivity is a matter of the extrinsicproperties; e.g., development off-proper microstructure for strongmagnetic domain pinning. And, this is generally not associated with asingle phase single crystal as suggested by the SEM and XRDexaminations, where FIG. 1 and FIGS. 2a-2c , did not reveal anyelemental precipitations, segregations, or/and any regular grainstructure/microstructure. The detailed high resolution TEM examinationof the as-grown sample of EXAMPLE #1, FIG. 7a , showed its basicuniformity and integrity with rare and small Co-enriched and Cu depletedregions/stripes which were coherent with the matrix. Unfortunately, thesize of these stripes did not allow clear EDS composition determinationor/and structural analysis. However, these may be embryonic structuraldefects caused by stacking faults trying to compensate various channeldisorders of the material, although the inventors do not wish to bebound by any theory in this regard. On the other hand, these also mightbe the nucleation set points for the spinodal decomposition, althoughagain the inventors do not wish to be bound by any theory in thisregard.

Although in EXAMPLE #1 the Ta content was incorporated unintentionally,the following heat treatment experiments demonstrate a role that Taplays a beneficial role in revealing and taking part in the formation ofnecessary microstructural modifications that improve magneticproperties; e.g., coercivity of the material.

The heat treatment procedure included dwelling at 1040° C. for 10 hoursand then cooling down at 10° C./minute to 400° C. for magnetichardening, dwelling at this temperature for next 8 hours with subsequentfurnace cool to room temperature. Phase analysis of powder X-ray patterntaken from crushed heat treated material, FIG. 4, revealed clearpresence of Ta-like impurities (Fm-3m, a=4.446(1) Å) with a majorityphase still retaining CaCu₅-type structure (P6/mmm, a=4.944(1),c=4.029(1) Å). SEM back scattered electron images taken from [001]surface of the heat treated crystals clearly contrasted to similar onesfrom the as-grown crystals, FIGS. 2a-2c , i.e., showing degradation ofthe uniformity and/or single phase nature of the crystal. Therefore the“single crystal” will no longer be used for accurate description of thematerial. Instead, the alternative tentative term “COMPOSITE CRYSTAL”will be used hereafter.

The “COMPOSITE CRYSTAL” exhibits bimodal microstructure that consists ofdarker matrix and lighter Ta-containing laminas in the single crystalmatrix. These laminas, seemingly, fill-up the regular extended defectsthat formed in the single grain single crystal during the heattreatment. These laminas strictly follow the hexagonal symmetry of theoriginal crystal, crossing each other at 60° and/or 120° angles ofintersection, FIG. 5a-5d . The thickness of the laminar features wasabout 0.05-0.1 μm, and their lengths varied in the range of about 1 toabout 10 μm. Distance (spacing) between two “parallel” laminas typicallywas about 2 to about 3 μm.

The elemental SEM X-ray analysis of the target areas 1-6 of the heattreated material appear in FIG. 6b and Tables 2 and 3, and clearlyindicates segregation of Ta and/or Ta-rich phase into the laminarfeatures, whereas only minor Ta concentrations are detected in thematrix material, Table 2. The TEM analysis confirmed that these are90-95% pure rectangular blocks of Ta (according to diffraction patternsand elemental analysis), and their interfaces were coherent with thematrix material, FIG. 7b .

TABLE 2 Atomic Fractions Area Co Cu Ce Ta 1 67.19 16.94 15.74 0.14 267.23 17.00 15.65 0.12 3 66.76 17.30 15.85 0.09

TABLE 3 Atomic Fractions Area Co Cu Ce Ta 4 62.20 14.96 14.97 7.87 559.41 13.88 14.69 12.0 6 66.66 17.21 15.66 0.47

Formation of such “COMPOSITE CRYSTAL” appears to be responsible forprofound change/improvement of magnetic properties after the heattreatment, i.e., significant increase of H_(c) from about 1.6 to about6.3 kOe yet with increase of B_(r) (M_(s)) from about 4.2 (4.2) to about5.3 (5.7) kG, resulting in (BH)_(max) of about 7.8 MGOe, FIGS. 8a-8c .Thermodynamic transformation of single crystals during the heattreatment appears to create a 3D array of extended defects withinprimarily single grain single crystal. Although not wishing or intendingto be bound by any theory, possible causes of this may be associatedwith Ta atoms leaving the matrix interstices at lower temperatures asindependent thermodynamic event and/or in coincidence with matrixdegradation induced by deceased miscibility at lower temperatures. Thesedefects then may serve as pinning sites for magnetic domains creatingnecessary extrinsic conditions for magnetic coercivity.

Example #1A

In an attempt to reproduce the results of EXAMPLE #1 but in bulk ingotform and also for scaling-up material preparation, an approximate 8 gramare-melted button (ingot) was prepared to have the Ta-doped alloycomposition of EXAMPLE #1; i.e., Ce_(15.5)Ta_(0.6)Co_(67.8)Cu_(16.1).The alloy was prepared by arc-melting elemental constituents on awater-cooled copper hearth under partial vacuum with purified argon,rotated/flipped and re-melted twice for the homogenization.

FIG. 9a shows magnetic properties of both the bulk, non-alignedpolycrystalline button (ingot) and the single crystal material after thesame heat treatment; i.e., 1040° C. (10 hours)→cooling 10° C./min.→400°C. (8 hours)→furnace cool to room temperature, i.e., cooling in theturned-off furnace to room temperature. Seemingly, both bulk and singlecrystal materials show nearly the same coercivity values and differ inB_(r) and M_(s), FIG. 9a . The latter result is anticipated consideringpolycrystalline and non-aligned nature of the bulk sample. FIG. 9b alsoreveals a remarkable approximate 390° C. Curie temperature, which wasevaluated by magnetic TGA method.

Example #2

This example illustrates that Fe-for-Co substitutions in the Ta-dopedcomposition of EXAMPLE #1 can result in improvement of saturationmagnetization up to about 60 to about 65%.

For example, an approximate 8 gram are-melted button (ingot) wasprepared to have a Fe-modified alloy composition; i.e.,Ce_(15.5)Ta_(0.6)Co_(57.6)Fe_(10.2)Cu_(16.1) by arc-melting byarc-melting elemental constituents on a water-cooled copper hearth underpartial vacuum with purified argon, rotated/flipped and re-melted twicefor the homogenization.

FIG. 10 shows comparison hysteresis loops of the Fe-substitutedcomposition bulk, unaligned sample and of the original bulk unalignedsample of EXAMPLE #1. FIG. 10 shows improvement of both coercivities andmagnetizations in the Fe-substituted sample; i.e., H_(c) from about 6.3to about 7.6 kOe, B_(r) from about 2.1 to about 4.7 kG and M_(s) fromabout 3.2 to 5.2 kG.

The present invention envisions that up to about 20 atomic % of Co canbe substituted by less significantly expensive Fe with improvement ofsaturation magnetization up to about 60 to about 65%.

Example #3

This example illustrates successful growth of single crystals of bothTa-, and Fe-doped Ce/Co/Cu permanent magnet, i.e.,Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4).

Similarly as in EXAMPLE #1, well-formed plate-like crystals wereobtained during self-flux single crystal growth from an initial loadingcomposition Ce₁₈Co_(55.8)Fe_(9.8)Cu_(16.4) in a Ta container at1050-1070° C. The excess of flux was decanted by centrifuging at nearlyreaction temperature. After performing the heat treatment identical toEXAMPLE #1, i.e., 1040° C. (10 hours)→cooling with rate 10° C./min.→400°C. (8 hours)→furnace cool to room temperature; i.e., cooling in theturned-off furnace to the room temperature, the EXAMPLE #3 showedsignificant improvement of magnetic energy characteristics reaching(BH)_(max.)=about 13 MGOe, FIGS. 11a , 11 b.

Cobalt content is decreased by about 6 at. % in comparison to EXAMPLE#1, while B_(r) increased by about 40% (to about 8 kG) in heat treatedFe-doped sample and H_(c) in the Fe-doped sample is about 60% (about 0.4T) from Fe-free sample, FIGS. 11a ,11 b.

FURTHER EXAMPLES

Table 4 represents compositions and main magnetic characteristics of theexperimental examples that are embodiments of the invention. Bothsingle-crystalline and polycrystalline synthetic approaches were usedfor sample preparation. Detailed description of the synthetic approachesis presented below.

In an attempt to reproduce the results of single-crystalline examples(see Table 4, EXAMPLES #3a, 4a and 6a) in bulk ingot form and also forscaling-up material preparation, an approximate 8 gram arc-meltedbuttons (ingots) were prepared to have the Ta-doped alloy composition ofEXAMPLE #3, 4 and 6, respectively. Also the polycrystalline arc-meltedexamples were tested with various Ta contents, i.e., EXAMPLES #7-9. Thealloys were prepared by arc-melting elemental constituents on awater-cooled copper hearth under partial vacuum with purified argon,rotated/flipped and re-melted twice for the homogenization.

The EXAMPLE #10 was synthesized in the alumina crucible, jacked in afused silica, under the argon gas atmosphere. This synthetic approachwas used to test reactivity of the components with the alumina crucible,since the alumina crucible syntheses are common casting techniques. Theexperiment confirmed that the Ce/Co/Cu gap magnets can be prepared inalumina crucibles. Magnetic characteristics of the EXAMPLE #10 (seeTable 4) are comparable to the single crystal growth results of EXAMPLE#3. The EXAMPLE #11 represents a Hf-doped polycrystalline sample (seeTable 4) prepared by arc-melting as described above for EXAMPLES 7-9.

Table 4 shows magnetic properties of all the bulk, wax-alignedpolycrystalline button (ingot) and the single-crystalline material afterthe same heat treatment; i.e., 1040° C. (10 hours)→cooling 10°C./hour→400° C. (8 hours)→furnace cool to room temperature, i.e.,cooling in the turned-off furnace to room temperature. Seemingly,prepared by different methods (single crystal growth, arc-melting,alumina crucible) and in small (3 gram) and larger quantities (8 gram),these systems reach comparable energy characteristics. This confirmshigh reproducibility of the initial experiments and shows the goodperspectives for scaling-up, manufacturing and mass production.

TABLE 4 Experimental examples of the Ce/Co/Cu gap magnets containingrefractory metal (Ta or Hf) and optionally Fe, with compositions,synthesis methods and magnetic characteristics. The best examples areemphasized by gray shading. The experimental examples 1a and 2 (seeabove) were not aligned, thus their magnetic characteristics are notfinalized and not presented in the Table. Example Synthesis Composition,at. % ρ M_(s) B_(r) H_(c) (BH)_(max.) T_(c) # method Ce Ta(Hf) Co Fe Cu(g/cm³) (kG) (kG) (kOe) (MGOe) (K) 4 Single crystal 15.1 1.0 74.4 — 9.58.5 6.7 5.9 0.4 1.0 670 growth I  4a Polycrystalline, 15.1 1.0 74.4 —9.5 8.5 5.8 5.5 1.5 3.8 n/a arc melted 5 Single crystal 16.3 0.6 68.9 —14.2 8.4 5.9 5.5 2.9 6.4 515 growth II 1 Single crystal 15.7 0.6 67.8 —15.9 8.5 5.8 5.5 6.3 7.8 490 growth III 6 Single crystal 16.3 0.3 61.7 —21.7 8.5 4.1 3.8 8.2 3.4 450 growth IV  6a Polycrystalline, 16.3 0.361.7 — 21.7 8.5 4.2 3.8 9.1 3.2 n/a arc-melting 3 Single crystal 14.31.0 62.0 12.3 10.4 8.3 8.1 8.0 3.4 12.8 820 growth V 7 Polycrystalline,15.4 0.1 62.0 12.0 10.5 8.4 6.6 6.2 4.4 7.5 n/a arc-melting 8Polycrystalline, 15.0 0.5 62.0 12.0 10.5 8.4 7.7 7.4 3.9 10.2 n/aarc-melting  3a Polycrystalline, 14.3 1.0 62.0 12.0 10.4 8.4 8.1 7.8 3.811.0 n/a arc-melting 9 Polycrystalline, 14.2 1.3 62.0 12.0 10.5 8.4 7.37.2 3.3 9.4 n/a arc-melting 10  Polycrystalline, 16.6 0.1 59.3 13.7 10.38.3 7.3 7.2 4.2 12.4 695 alumina crucible 11  Polycrystalline, 15.0 0.761.8 12.1 10.4 8.4 7.6 7.3 5.2 12.5 ~800  arc-melting (Hf)

Note that EXAMPLE #11 representing the Hf-doped polycrystalline sample(see Table 4) showed the best combination of magnetization andcoercivity among all examples presented in Table 4; e.g. see FIG. 24.Example #11 showed a microstructure having Hf-containing laminas in thematrix in similar fashion as the Ta-doped samples having Ta-containinglaminas after the above heat treatment (see FIG. 23a -23b and FIG. 23ccontaining Table 8.

Single-Crystalline Samples

The following examples demonstrate synthesis, structure, and magneticproperties of Ta-, Cu- and Fe-substituted CeCo₅ magnet alloys. Using aself-flux technique, single crystals of sampleI—Ce_(15.1)Ta_(1.0)Co_(74.4)Cu_(9.5), sampleII—Ce_(16.3)Ta_(0.6)Co_(68.9)Cu_(14.2)), sampleIII—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9), sampleIV—Ce_(16.3)Ta_(0.3)Co_(61.7)Cu_(21.7), (EXAMPLE and sampleV—Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4) were grown. The singlecrystals III and V correspond to EXAMPLE #1 and EXAMPLE #3 which werementioned earlier as initial/provisional results (see above). Acomprehensive and detailed characterization of the samples is presentedbelow with respect to magnetic behavior and unique magneticcharacteristics; i.e., coercivity mechanism.

EXPERIMENTAL

Single crystals were grown via the solution growth method described byP. C. Canfield et al. “Growth of single crystals from metallic fluxed”,Philos. Mag., 65, 1117-1123 (1992) and P. C. Canfield et al. “Propertiesand Applications of Complex Intermetallics, Solution Growth ofIntermetallic Single Crystals: A Beginner Guide”, edited by Belin-Ferre,Chap. 2, (World Scientific, Singapore 2010), the teachings of which areincorporated herein by reference to this end. The reaction metals (Ce(99.99%), Cu (99.95%) from Ames Laboratory MPC (Material PreparationCenter) and Co (99.95%) from Alfa Aesar) were placed into 3-capped Tacontainers (see reference 40) welded under an Ar atmosphere, which thenwere sealed into fused silica tubes and placed into a high-temperaturebox furnace. The furnace was heated from near room temperature to 900°C. over 3 hours, held at 900° C. for 3 hours, heateds to 1200° C. overthree more hours, and held at 1200° C. for 10 hours. The furnace wasthen cooled to 1070° C. over 75 hours. At 1070° C. the excess flux wasdecanted by centrifuging (see Canfield references above). Decanting tookplace as the centrifuge accelerated from rest toward a 8.5 krpm setpoint.

Heat Treatment:

After growth, some single crystals underwent identical, two-stage, heattreatments performed in a Dentsply Ceramico (Vulcan 3-Series)multi-stage programmable furnace, which included dwelling at 1040° C.for 10 hours, then cooling at a rate of 10° C./min to 400° C. followedby dwelling at this temperature for the next eight hours with asubsequent furnace cool to room temperature. Different Cu contents mayrequire slightly different temperature/time parameters for the bestfinal magnetic characteristics and can be determined empirically.

TABLE 5 Composition of single crystals (with standard deviation) andtheir lattice parameters as-grown and after the heat treatment.EDS-composition, at. % Lattice parameters Ce Ta*** Co Fe Cu a, c, 

 ; V, 

 ³ **** # ag* ht** ag ht Ag Ht ag Ht Ag ht Ag Ht I 15.1(1) 16.1(1) 1.00.6 74.4(2) 73.6(2) — — 9.5(1) 9.8(1) 4.912(1) 4.921(1) 4.045(1)4.031(1) 84.52(1) 84.58(2) II 16.3(1) 16.2(1) 0.6 0.4 68.9(2) 69.4(2) —— 14.2(1) 14.0(1) 4.933(1) 4.933(1) 4.031(1) 4.028(1) 84.95(2) 84.90(2)III 15.7(1) 15.8(1) 0.6 0.1 67.8(2) 67.1(2) — — 15.9(1) 17.1(1) 4.943(1)4.944(1) 4.032(1) 4.028(1) 85.31(1) 85.26(1) IV 16.3(1) 16.5(1) 0.3 0.0561.7(2) 61.9(2) — — 21.7(1) 21.6(1) 4.950(1) 4.95 

 (1) 4.033(1) 4.028(1) 85.57(2) 85.61(2) V 14.0(1) 13.9(1) 1.0 0.262.0(2) 62.7(2) 12.3(1) 13.0(1) 10.4(1) 10.2(1) 4.922(1) 4.924(1)4.075(1) 4.071(1) 85.50(2) 85.48(2) *aε-grown, **heat-treated: 10 

 0 C. (10 h) > [10° C./min]→ 400° C. (8 h), ***standard deviations forTa vary within 0.02-0.05 at. %; **** space group: P 6/mmm

Samples for metallographic examination were placed in 1 inch diameterepoxy resin pucks, and polished with various grits of silicon carbidefollowed by a glycol-based, fine, polycrystalline, diamond suspension.Plate-like single crystals were mounted with their plates parallel tothe polishing surface to allow for characterization along planesperpendicular to the crystals [001] direction. Imaging studies of singlecrystal samples were per-formed on an FEI Teneo field emission scanningelectron microscope. Their compositions were determined via energydispersive X-ray spectra obtained using an Oxford EDS/EBSD moduleaveraging over 3-5 regions on their metallographically prepared surfaces[see Table 5].

Tem Characterization:

Cross sections from single crystal sample III were prepared using adual-beam focused ion beam system (FEI Helios NanoLab G3 UC) with alift-out approach. To reduce surface damage sustained during Ga ionmilling, the final thinning and cleaning step were conducted using 5 kVand 2 kV for 5 min. The TEM analysis was performed on a Titan Themis(FEI) probe Cs-corrected TEM equipped with a Super-X EDS detector tocharacterize microstructure and elemental distribution.

Powder and Single Crystal X-Ray Diffraction:

Polycrystalline powders were obtained by crushing the sample with anagate mortar and pestle. X-ray power diffraction data were collectedfrom the as-grown and heat-treated crystals. The measurements wereper-formed using PANalytical X-Pert Pro (Co K_(α)—radiation, λ=1.78897Å) and Bruker D8 Advance (Cu K_(α)—radiation, λ=1.54056 Å) diffractionsystems. Powdered samples were evenly dispersed on a zero-background Si—holder with the aid of a small quantity of vacuum grease. Diffractionscans were taken in the θ/2θ mode with the following parameters: 2θregion: 20-110°, step scan: 0.02°, counting time per step: 60 s. TheFullProf Suite program package (see reference 41) was used for Rietveldrefinement of the crystal structures.

Single-crystal diffraction data were collected at room temperature usinga Bruker SMART APEX II diffractometer (Mo K_(α)—radiation) equipped witha CCD area detector. Four sets of 360 frames with 0.5° scans in ω andexposure times of 10-15 s per frame were collected. The reflectionintensities were integrated using the SAINT subprogram in the SMARTsoftware package, Bruker AXS Inc., Madison, Wis. 1996. The space groupwas determined using the XPREP program and the SHELXTL 6.1 softwarepackage, Bruker AXS Inc., Madison, Wis. 2000. Empirical absorptioncorrections were made using the SADABS program (R. H. Blessing, “Anempirical correction for adsorption”, Acta.Crystallographica Section AFoundations of Crystallogrpahy, 51, 33-38 (1995). Finally, eachstructure was solved by direct methods using SHELXTL 6.1 and refined byfull-matrix least-squares on F₀ ², with anisotropic thermal parametersand a secondary extinction parameter.

Magnetic Properties Measurements:

Magnetic property measurements were obtained using a vibrating samplemagnetometer in a cryogen-free VersaLab physical property measurementsystem (Quantum Design) with magnetic fields up to 3 T and temperaturesin the 50-350 K range using the standard option and 300-1000 K rangeusing the oven option. An alumina cement (Zircar) was used to hold thesample on the heater stick for the high-temperature measurements. Thedemagnetization factors are determined experimentally using the relationH_(int.)=H−NM.

Structure and Composition Analysis:

SEM/EDS Examinations and Composition Analysis:

The SEM backscattered electron images of the as-grown crystals [FIG.13a-13c , upper panels] display the uniformity of their (0001) polishedsurface (even at ×30,000 magnification) which suggests a single-phase.Elemental EDS analysis [Table I] showed the Ce,Co/Cu ratios are close tothe 1:5 stoichiometry with Cu contents increasing from ˜10 to ˜22 at. %,corresponding to 12-26% of Co/Cu substitution. With respect to Cecontent, crystal sample I and sample III contain 15-15.7 at. %, which islower than the Ce content in sample II and sample IV and significantlylower than 16.7 at. % Ce content required for the exact 1:5 typestoichiometry. Also a minor presence of Ta (0.3-1 at. %) was detected inall five samples. The Ta content appears to be correlated to the Cucontent as seen in [Table 5]. The presence of Ta is explained by theslight dissolution of the inner walls of the Ta reaction container anddiffusion of Ta atoms into the liquid at high temperatures. Since no Taprecipitation or segregation was observed in the SEM/EDS analysis of theas-grown crystals, Ta is either being incorporated into the crystalstructure as interstices or as uniformly distributed nano-scaleprecipitates, although the inventors do not intend or wish to be boundby any theory in this regard. However the slight Ce depletion and thepresence of Ta suggest the possibility of minor deviations from theclassic CaCu₅-type crystal structure towards various channel disordersor “dumb-bell” problems characterized elsewhere by O. Bodak et al.,“Structural and magnetic properties of iron-rich compounds in theYb—Fe—Al system”, Journal of Alloys and Compounds 354, L10-L15 (2003);Ya. O. Tokaychuk et al., “Structural and magnetic properties ofiron-rich compounds in the Yb—Fe—Ga system”, Journal of Alloys andCompounds 415, 8-11 (2006); and Radovan Cerny et al., Local atomic orderin the vicinity of Cu2 dumbbells in TbCu₇-type YCu_(6.576) studied usingbragg and total scattering techniques”, Intermetallics 17, 818-825(2009.

The SEM back scattered electron images taken from the (0001) surface ofthe heat treated crystals, [FIG. 13a -FIG. 13c lower panels], showdegradation of the single phase crystal into a bimodal microstructureconsisting of a darker matrix and lighter laminas. These laminas followthe hexagonal symmetry of the original crystal crossing each other at60° or 120° angles. The thickness of the laminar features is ˜0.05-0.1μm, and their lengths vary in the range ˜1-10 μm. Distances between twolaminas are ˜2-3 μm. The elemental EDS analysis of the heat treatedmaterial [Table 5] indicates the segregation of Ta-rich phases into thelaminar features, whereas the matrix material becomes practicallyTa-free in the Cu-richest crystal sample IV.

X-Ray Crystal Structure Determination:

Powder and single crystal X-ray analyses were performed to determine thestructure of crystals samples I-V. Rietveld fitting of the powder X-raypattern taken from the as-grown, crushed and thoroughly powdered, singlecrystals of samples I-V showed that all Bragg reflections were wellindexed within the CaCu₅-type structure (hP6, P6/mmm), providing strongevidence for the single-phase nature of the as-grown crystals inagreement with our SEM observations [FIG. 13a-13c ]. To address theEDS-observed Ta presence and Ce deficiency, especially in the as-growncrystals samples I, III, and V [see Table 5], known structuralderivatives of CaCu₅ were considered (O. Bodak et al., “Structural andmagnetic properties of iron-rich compounds in the Yb—Fe—Al system”,Journal of Alloys and Compounds 354, L10-L15 (2003)). These derivativesare typically observed in binary and ternary rare-earth—transition metalsystems near the 1:5 and 2:17 stoichiometries and result from thereplacement of rare-earth atoms by pairs of transition metal atoms. TheCaCu₅ substructure can be retained if the replacement is fully random,as in TbCu₇, but may be transformed into various superstructures, suchas Th₂Zn₁₇, Th₂Ni₁₇, etc., if the substitution is ordered. A thirdpossibility comes as combination of ordered and disordered rare-earth;i.e. “dumb-bell” substitutions which are contained in a superstructure,e.g., LuFe_(9.5) and PrFe₇. Rietveld refinements with structural modelsallowing the presence of Ta were tried but the clear indexing of Braggreflections within the parent, CaCu₅-type, 1:5 structure indicated aminor and random distribution of Ta.

Single crystal structure solutions of samples III-V confirmed theirCaCu₅ substructure (Tables 6, 7). However, disorder was detected withinthe 1D hexagonal channels, i.e., the residual electron density peaks ofabout 5.0, about 3.8 and about 13.2 e/Å³ at (0 0 z), z about 0.295 forsamples III, IV and V, respectively. Only by filling the 2e site withthe heaviest and largest available pair, Ta-Ta, was it possible to reachsatisfactory refinement. The R1/wR2 residuals dropped by 50-70% incomparison to the solutions without Ta and showed minimal fluctuationsof the rest electron density in the final fits. Differential Fouriermaps for samples III-V with and without the “dumb-bell” showed disorder.One significant deficiency of the solutions is the interatomic T-Tdistances of 2.35 Å, which is typical for Co—Co, Co—Cu and Co about Fepairs but is extremely short for Ta-Ta. However, the “dumb-bell”configuration with large and heavy atoms similar to Ta is notunprecedented and was reported for similar structure ofCeFe₁₀Zr_(0.8)(d_((Zr-Zr)) 2.65 Å). However, the stoichiometry of sampleV shows significant deviation from the ideal 1:5 stoichiometry. Thecontent of 1D channels (Ce plus the Ta-Ta pairs) does not reach theexpected 16.7 at. %, meaning that some of Ta atoms must participate inthe channel disorder, although the inventors do not wish or intend to bebound by any theory in this regard.

TABLE 6 Single crystal and Refinement Data for III -Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9), IV -Ce_(16.3)Ta_(0.3)Co_(61.7)Cu_(21.7), V -Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4). Crystal III IV V EDScomposition Ce_(0.94)Ta_(0.04)Co_(4.06)Cu_(0.94)Ce_(0.99)Ta_(0.00)Co_(3.70)Cu_(1.30)Ce_(0.86)Ta_(0.06)Co_(3.72)Fe_(0.73)Cu_(0.62) refined compositionCe_(0.98)Ta_(0.04)Co_(4.25)Cu_(0.75)Ce_(0.99)Ta_(0.02)Co_(3.79)Cu_(1.21)Ce_(0.86)Ta_(0.12)Co_(3.68)Fe_(0.72)Cu_(0.60) formula mass 442.68 442.57449.52 space group; Z P6/mmm; 1 P6/mmm; 1 P6/mmm; 1 a (Å) 4.946(1)4.952(1) 4.928(1) c (Å) 4.038(1) 4.035(1) 4.073(1) V (Å³) 85.57(4)85.70(5) 85.66(2) d_(c) (Mg/m³) 8.52 8.57 8.69 μ (mm⁻¹; Mo Kα) abs coef37.85 37.08 39.78 reflns collected/R_(int)  1631/0.025  2002/0.042 1808/0.027 ind. data/restrains/params 79/0/12 109/0/13 91/0/11 GoF (F²)1.221 1.129 1.172 R1/wR2 [I > 2σ(I)] 0.018/0.041 0.021/0.048 0.030/0.063R1/wR2 [all data] 0.021/0.041 0.025/0.046 0.031/0.063 Largest diffpeak/hole (e/Å³)  0.80/−0.74  1.04/−0.99  1.91/−1.53

TABLE 7 Atomic coordinates, Equivalent Isotropic Displacement Parameters(Å × 103), and Site Occupancy Factors Refined for III -Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9), IV -Ce_(16.3)Ta_(0.3)Co_(61.7)Cu_(21.7), V -Ce_(14.3)Ta_(1.0)Co_(62.0)Fe_(12.3)Cu_(10.4). atom WP X Y z U_(eq) SOF #Ce 1a 0 0 0 15(1) 0.977(2) III 16(1) 0.988(2) IV 19(1) 0.936(3) V Ta 2e0 0 0.280(6) 15(1) 0.023(2) III 0.296(9) 16(1) 0.012(1) IV 0.292(4)19(1) 0.064(3) V M1^(a) 2c ⅔ ⅓ 0 14(1) 1.00 Co  III 15(1)  0.23(7) Cu IV23(1) 1.00 Co  V M2 3g ½ 0 ½ 10(1)  0.25(6) Cu III 10(1)  0.25(5) Cu IV10(1) 0.24 Fe/0.20 Cu V ^(a)The atomic symbol “M” stands for Co/Cu orCo/Fe/Cu mixed occupancy; 3g occupancy for sample V have been fixed.

FIG. 14a is a high angle annular backfield (HAADF) scanning transmissionelectron microscope (STEM) image of as-grown sampleIII—Ce_(15.7)Ta_(0.6)Co_(67.8)Cu_(15.9) showing the overallmicrostructure. The entire region consists of a single crystallinephase. FIG. 14b is an enlarged HAADF image which shows a dark-contrastline, which was the only feature which could be found in the entire scanarea. FIG. 14c is a diffraction pattern taken from the region shown inFIG. 14b including the matrix and the dark line. It clearly shows thesingle crystalline 1:5 phase. It seems that the dark line region has thesame crystal structure and it is not a precipitate which would have madeadditional diffraction spots in FIG. 14c . FIG. 14d is a high resolutionSTEM image taken from left boxed area in FIG. 14b under [1-10] zoneaxis. The bright spots and the dark spots in the images correspond toatomic columns of Ce and Co/Cu elements, respectively. The inset atbottom right is an enlarged atomic image with atomic model of hexagonal1:5 Ce/Co/Cu phase. The bright dots and dark dots in the imagescorrespond to atomic columns of Ce and (Co, Cu) elements, respectively.FIG. 14e is an enlarged image of right boxed area in FIG. 14b and thedark line in the single crystalline phase is shown clearly. FIG. 14fshows EDS elemental mapping of FIG. 14e . The chemical contrast betweenthe matrix and the dark line is present. The EDS result shows the darkline is Co-enriched and Cu deficient. The small Co and Cu elemental maps(insets in FIG. 14f Co are presented for contrasting observation of Cudepletion in the same dark line.

FIG. 15a is a HAADF image of an annealed sample showing the overallmicrostructure. Many bright areas were observed unlike the un-annealedsample shown before in FIG. 14a-14b . FIG. 15b is an enlarged image ofthe circled area in FIG. 15a . FIG. 15c is the EDS elemental mappingcorresponding to FIG. 15b . The bright regions in FIG. 15b are Ta-richand considered as Ta precipitates. Additionally, a few dark lines areobserved in the Ta precipitate. The difference in brightness ofprecipitates is attributed to the difference in the thickness of eachprecipitate. FIG. 15d is an enlarged image of the boxed area in FIG. 15b, and FIG. 15e is a diffraction pattern taken from FIG. 15d includingthe matrix and the Ta-containing precipitates. FIG. 15d shows Taprecipitates coherently embedded by epitaxial precipitation and thecorresponding diffraction pattern shows the epitaxial relationshipbetween the matrix and Ta precipitate. The orientation relation wasobserved as follows: (110) Ce—CoCu//(110) Ta; (002) CeCoCu//(1-10) Ta;and [1-10] CeCoCu//[001] Ta. FIG. 15f is a high resolution STEM imagetaken from the matrix in FIG. 15d under [1-10] zone axis. It is the sameas that seen in FIG. 14d . The inset at bottom right is an enlargedatomic image with an atomic model of hexagonal 1:5 Ce/Co/Cu phase. Thebright dots and dark dots in the images correspond to atomic columns ofCe and Co/Cu elements, respectively.

FIGS. 16a-16f show high resolution HAADF images of the interface betweenthe matrix and the Ta precipitate taken from right-hand-boxed area inFIG. 15b and corresponding EDS elemental mapping results [FIG. 16b-16f]. The white dashed lines indicate the same position in each image.Although Cu-rich and Co-deficient region was observed near theprecipitate, there was also Co, Ce-rich and Cu-deficient interfacebetween the matrix and the Ta precipitate. The dark lines in the Taprecipitate turned out to be Co-rich. Considering EDS maps at theinterface and near the precipitate, Co possibly infiltrated into theprecipitate [FIG. 15d ], and Co became deficient near the precipitatewith relative Cu-enrichment as a result, although the inventors do notintend or wish to be bound by any theory in this regard.

Magnetic Properties:

Curie temperature, magnetocrystalline anisotropy field, and energydensity of as-grown crystals: FIG. 17 presents the Curie temperaturesfor samples I-V as inferred from the peak in dM/dT shown in the inset.The Curie temperatures T_(c) estimated by the minimum in the derivativecorrespond closely to the T_(c) derived via the more accurate Arrot plotmethod (see below). The T_(c)-value decreases rapidly with increasing Cucontent for Fe-free samples I-IV. This indicates weakening in theferromagnetic exchange interactions within the Co sublattice due to theintroduction of nonmagnetic Cu. In contrast, the Fe-doped crystal sampleV shows remarkable improvement of T_(c), increasing by over 150 K to 820K, a value that is significantly higher than the T_(c)=653 K of theparent CeCo₅. Band structure analysis indicates that Fe— doping of CeCo₅and Ce(Co,Cu)₅ increases the ordering energy ΔE=E_(NM) E_(F M)(NM=non-magnetic and FM=ferromagnetic states), as well as the totalmagnetic moment of the systems. This leads to the remarkable increase ofthe Curie temperature and saturation magnetization.

To more formally determine T_(c), an Arrot plot analysis was conductedfor sample III using isotherms between 460 K and 500 K. The Curietemperature for sample III was estimated to be 480 K, since the isothermat that temperature was closest to a straight line and passes throughthe origin. FIG. 17 shows representative M(H) for each H=0. As can beseen, these data suggest a T_(c) 820 K (estimated by generalized Blochlaw fitting of spontaneous magnetization), in good agreement with FIG.16.

The magnetocrystalline anisotropy field, H_(a), at room temperature wasdetermined for all as-grown crystal samples I-V. The low temperatureestimations of H, were conducted for crystal sample III and V. Theanisotropy field was estimated by the high-field, linear extrapolationof the filed-dependent moment along the easy axis [001] and hard (H⊥[001]) axis (see E. A. Nesbitt et al. “Intrinsic magnetic properties andmechanism of magnetization of Co—Fe—Cu—R permanent magnets,” in AIPConference Proceedings (AIP 1973) and Tej N. Lamichhane et al.,“Ce_(3-x)Mg_(x)Co₉: Transformation of a Pauli paramagnet into a strongpermanent magnet,” Physical Review Applied 9 (2018)).

The room temperature H_(a) for the Fe-free, as-grown crystal samplesI-IV exhibit a maximum anisotropy field of about 118 kOe (in crystalsample II). The addition of Fe showed a detrimental influence on themagnetocrystalline anisotropy, (e.g. in Fe-doped sample V, theanisotropy field dropped to about 65 kOe), but the spontaneousmagnetization increased by about 30% compared to crystal samples withsimilar Cu contents. Low temperature measurements estimate thespontaneous magnetization for crystal samples III and V to be about 3.7and about 6.8 μ_(B)/f.u., respectively.

The temperature dependent magnetocrystalline anisotropy energy densitywas measured using the Sucksmith-Thompson method by using the hard axismagnetization iostherms for crystal samples III and V magnetizationisotherms for crystals samples III and V [FIG. 19]. For a description ofthe Sucksmith-Thompson method, see Tej N. Lamichhane et al.,“Ce_(3-x)Mg_(x)Co₉: Transformation of a pauli paramagnet into a strongpermanent magnet,” Physical Review Applied 9 (2018); W. Sucksmith and J.E. Thompson, “The magnetic anisotropy of cobalt,” Proceedings of theRoyal Society of London A: Mathematical, Physical and EngineeringSciences 225, 362-375 (1954), and Valentin Taufour et al., “Structuraland Ferromagnetic Properties of an Orthorhombic Phase of MnBi Stabilizedwith Rh Additions,” Phys. Rev. Applied 4, 014021 (2015).

Interestingly, the as-grown crystals showed magnetic hysteresis whenmeasured along the easy axis of magnetization [001]. For example,crystal sample III exhibited a hysteresis, which reached H_(c)≈1.6 kOeand B_(r)≈4.2 kG, M_(s)≈4.2 kG and (BH)_(max.)≈3.5 MGOe [FIG. 20], whichis comparable to most of the anisotropic sintered alnico grades. This isremarkable considering the common belief that the appearance of thecoercivity is a result of the extrinsic properties, e.g., development ofproper microstructure for strong magnetic domain pinning, and this isgenerally not associated with a single phase single crystal asdetermined by the SEM examinations [FIGS. 13a-13b ] and XRDexaminations, which did not reveal any elemental precipitations,segregations, or any microstructure on their correspondinglength-scales.

The detailed high resolution STEM examination of the as-grown sample III[FIG. 14a-14c ] showed the basic uniformity and integrity with smallCo-enriched and Cu depleted regions/stripes coherently dispersedthroughout the matrix. Unfortunately, the size of these stripes did notallow for EDS composition determination or structural analysis. Thesestripes may be embryonic structural defects caused by stacking faultscompensating for various channel disorders within the material and/orthe nucleation sites for the decomposition and/or miscibility gap,although the inventors do not intend or wish to be bound by any theoryin this regard.

Heat Treated Crystals—Coercivity, Pinning, and Magnetic Energy:

After heat treatment, crystal samples I-V showed significantly increasedmagnetic hystereses with a monotonic increase of coercivity, H_(c), anda linear decrease of spontaneous magnetization M_(s) with increasing Cucontent [FIG. 21]. For example, the magnetic characteristics of sampleIII change as follows: significant increase of H_(c) from ˜1.6 to ˜6.3kOe with an increase of B_(r) (M_(s)) from ˜4.2 (4.2) to ˜5.3 (5.7) kG,resulting in (BH)_(max.) of ˜7.8 MGOe [FIG. 22]. In addition to theconspicuous increase in magnetic hysteresis, there is a noteworthyincrease in saturated magnetization of the heat-treated samples [FIG.21, upper inset].

Referring back to EXAMPLE #3 (crystal V) to this same end, EXAMPLE 3#showed significant improvement of magnetic energy characteristicsreaching (BH)_(max.) of about 13 MGOe, FIGS. 11a, 11b . Fe thus stronglyimproved both the Curie temperature (to about 820 K) and themagnetization (to about 8 kG) of the heat-treated magnetic materialresulting in magnetic energy of about 13 MGOc at room temperature.

The increases in magnetic properties after the heat treatment correlatewith the appearance of the Ta-rich precipitates [see SEM images above,FIG. 13a-13c ]. The STEM analysis confirmed that these are 90-95% purerectangular blocks of Ta (according to diffraction patterns andelemental analysis), and their interfaces were coherent with the matrixmaterial. However, high magnification TEM EDS maps [FIG. 14c , FIG.16a-16f ] observed a Cu-deficient and Co-enriched layer at the interfaceof the precipitates and the matrix, and Co was detected in precipitatesas lines, which somewhat resemble observations of rare Co-enriched andCu-depleted lines in the as-grown STEM examination [FIG. 14a-14f ].

These results suggest that the high coercivity may be explained by theTa-rich precipitates serving as pinning sites and can be described usinga simple domain pinning model. Typically, the coercive force isinversely proportional to the saturation magnetization for a particularmagnetocrystalline energy (H_(c)=AK/M_(s)l, where A—exchange constant,K—magnetocrystalline anisotropy, M_(s)—saturation magnetization andl—the distance between the precipitates. According to the equation, byincreasing the amount of pinning precipitates the volume fraction of thematrix material and magnetization M_(s) of the system is decreased. Alsothe distances l between the precipitates become shorter. As a result,the coercivity H_(c) increases. Thus, the H_(c) of the sample crystalsshould be directly proportional to the Ta content. However, theinventors observed the inverse proportionality: total Ta contentmonotonically decreases in crystals samples I through IV [Table 1],whereas the coercivity monotonically increased [FIG. 20].

In contrast, the H_(c) increase correlated directly with increasing Cucontent [FIG. 21, see both insets], also following the proposedprecipitation coercivity mechanism (see equation above). Pinning ofmagnetic domains should occur on the precipitates, the amount of whichis regulated by Cu, rather than Ta content. However, precipitates thatare clearly associated with Cu were not observed in the samples, exceptCu-depleted regions observed in STEM experiments [FIG. 14c , FIG.16a-16f ].

Although the inventors do not wish or intend to be bound by any theory,the Ta-rich laminar precipitates therefore may be considered as asecondary effect, which are believed to decorate the extended 3D defectsand structural imperfections that originate from Cu depleted and Coenriched lines observed in the as-grown crystals [FIG. 16a-16f ] andconsequently develop into the regions between Ta-rich precipitates andmatrix in the thermally aged crystals [FIG. 14c , FIG. 16a-16f ]. Theamount of these imperfections appear to increase with increasing Cucontent and lead to increased coercivity.

The examples described above thus demonstrate the synthesis of fivedifferent single crystals of Ta, Cu and/or Fe substituted CeCo₅ usingthe self-flux technique. The results can be summarized as Indicatingthat the crystals so produced retained a CaCu₅ substructure andincorporate small amounts of Ta in the form of “dumb-bells” filling the2e crystallographic sites within the 1D hexagonal channel with the 1a Cesite, whereas Co, Cu and Fe are statistically distributed among the 2cand 3g crystallographic sites. The as-grown crystals appeared to besingle phased and homogenous in composition. Their single crystallinityis confirmed by XRD, SEM and TEM experiments. However they also exhibitsignificant magnetic coercivities, which are comparable to mostanisotropic sintered alnico grades. After the heat treatment(hardening), magnetic characteristics significantly improve. Ta atomsappear to leave the matrix interstices of the as-grown crystals andprecipitate in the form of coherent laminas creating the so-called“COMPOSITE CRYSTAL”. The “COMPOSITE CRYSTAL”, formed during the heattreatment, appears to contain a 3D array of structural defects within aprimarily single grain single crystal, although the inventors do notwish or intend to be bound by the proposed explanation above.

To this same end, the mechanism of coercivity appears to be regulated byCu, and pinning occurs on the extended 3D defects and structuralimperfections that originate from Cu depleted and Co enriched linesobserved in the as-grown crystals and consequently develop into theregions between Ta-rich precipitates and matrix in the thermally agedcrystals. The structural defects form as a result of a thermodynamictransformation of the matrix material associated with its partialdecomposition and/or decreased miscibility during hardening process.Significant improvement of magnetization in the heat-treated samples maybe associated either with the transformation of the matrix phase or withthe removal of Ta from the matrix. Fe strongly improves both the Curietemperature and magnetization of the system, which is associated with astrong increase in the magnetic ordering energy. The peculiarthermodynamic transformations, which lead to intragranular pinning and aunique coercivity mechanism that does not require the typical processingfor the development of extrinsic magnetic properties, could be used tocreate permanent mag-nets with lowered processing costs. Furthercomposition—temperature—time optimizations may result in a criticalmaterial free and cost-efficient gap magnet with energy product above 7to about 15 MGOe and even up to about 16.5 MGOe.

The present invention is advantageous to provide for substitution of Smfully by less expensive Ce in a 1:5-type magnet together with smalladditions of Ta to provide magnetic characteristics suitable for the“GAP MAGNET” at significantly lower material costs. Moreover, use ofgrain development techniques is not strongly required for development ofsignificant coercivities, making permanent magnets pursuant to thepresent invention also a process efficient material.

Although the present invention has been described with respect tocertain illustrative embodiments, those skilled in the art willappreciate that the invention is not limited to these embodiments andthat changes and modifications can be made therein within the scope ofthe invention as set forth in the appended claims.

We claim:
 1. A permanent magnet alloy, comprising Ce, Co, Cu, and arefractory metal comprising at least one of Ta, Hf, Zr, Nb, Mo, and W.2. The alloy of claim 1 wherein the refractory metal comprises Ta or Hf.3. The alloy of claim 1 wherein the refractory metal comprises Ta andHf.
 4. The alloy of claim 1 wherein the refractory metal is present inan individual or collective amount of at least 0.1 atomic % up to about3 atomic %.
 5. The alloy of claim 1 wherein Fe is substituted for someof the Co.
 6. The alloy of claim 1 which has a CeCo₅ crystal structure.7. The alloy of claim 1 which is heat treated to develop a bimodalmicrostructure having refractory metal-containing laminas in a primarilysingle crystal or polycrystalline matrix.
 8. The alloy of claim 7 havingCu-enriched and Co-depleted near respective laminas.
 9. A permanentmagnet alloy, comprising, in atomic %, about 13 to about 17% Ce, about38 to about 70% Co, about 10% to about 30 atomic % Cu, optionally about10% to about 20% Fe, and comprising at least one of Ta, Hf, Zr, Nb, Mo,and W in an individual or collective amount greater than 0.1 up to about3% atomic %.
 10. The alloy of claim 9 which is heat treated to develop abimodal microstructure having refractory metal-containing laminas in aprimarily single crystal or polycrystalline matrix.
 11. The alloy ofclaim 10 having Cu-enriched and Co-depleted regions near respectivelaminas.
 12. The alloy of claim 10 which is subjected to a solutiontemperature followed by aging to develop the bimodal microstructure. 13.The alloy of claim 10 which exhibits magnetic characteristics acceptablefor a “GAP MAGNET”.
 14. The alloy of claim 13 wherein the heat treatedalloy exhibits a T_(c)>300° C., H_(c)=0.5-1.0 T, B_(r)=4-8 kG and(BH)_(max.)=7-15 MGOe.
 15. A heat treated permanent magnet, comprisingCe, Co, Cu, and refractory metal comprising at least one of Ta, Hf, Zr,Nb, Mo, and W, the magnet having a bimodal microstructure withrefractory metal-containing laminas disposed within a microstructurematrix.
 16. The magnet of claim 15 having Cu-deficient interface betweenthe matrix and the respective laminas.
 17. The magnet of claim 15wherein the laminas reside in a primarily single crystal matrix or apolycrystalline matrix.
 18. The magnet of claim 15 comprising, in atomic%, about 13 to about 17% Ce, about 38 to about 70% Co, about 10% toabout 30 atomic % Cu, optionally about 10% to about 20% Fe, and one ormore refractory metals in an individual or collective amount greaterthan 0.1 up to about 3 atomic %.
 19. The magnet of claim 15 which issubjected to a solution temperature followed by aging to develop thebimodal microstructure.
 20. The magnet of claim 15 wherein the heattreated alloy exhibits a T_(c)>300° C., H_(c)=0.5-1.0 T, B_(r)=4-8 kGand (BH)_(max.)=7-15 MGOe.
 21. A method of making a permanent magnet,comprising heat treating a permanent magnet comprising Ce, Co, Cu, and arefractory metal comprising at least one of Ta, Hf, Zr, Nb, Mo, and W todevelop a bimodal microstructure having refractory metal-containinglaminas in a matrix microstructure.
 22. The method of claim 21 whereinthe magnet comprises, in atomic %, about 13 to about 17% Ce, about 38 toabout 70% Co, about 10% to about 30 atomic % Cu, optionally about 10% toabout 20% Fe, and one or more refractory metals in an individual orcollective amount greater than 0.1 up to about 3 atomic %.
 23. Themethod of claim 22 wherein heat treating involves heating the magnet toa solution temperature followed by aging to develop the bimodalmicrostructure.
 24. The method of claim 23 wherein the solutiontemperature is 1000-1100° C. for a time followed by aging at 400-450° C.for a time to develop the bimodal microstructure.